Ceramic powder and dense insulating coating

Abstract

The application relates to the technical field of ceramic coating materials, and particularly discloses a ceramic powder and a dense insulating coating. The ceramic powder is a core-shell structure particle, the core-shell structure particle is composed of a core material and a wall material wrapped outside the core material, the core material comprises a lanthanum magnesium hexaaluminate phase and a skeleton adjusting phase, and the wall material comprises a magnesium aluminum spinel phase, a free alpha-Al2O3 phase and a filling component. The skeleton adjusting phase and the filling component are optimized, and the overall proportion is optimized, so that the coating formed after the ceramic powder is sprayed has low connected porosity and a low coating thickness change rate, thereby improving the ability of the ceramic coating to maintain insulation after being kept at 1100 DEG C.

Description

Technical Field

[0001] This application relates to the field of ceramic coating materials technology, and in particular to a ceramic powder and a dense insulating coating. Background Technology

[0002] The hot-end components of aero-engines are typically surrounded by functional components such as ignition, temperature measurement, and condition monitoring. In addition to heat insulation, the surfaces of high-temperature metal components also require maintaining local electrical insulation. Thermal spraying of ceramic coatings can form a high-temperature resistant insulating layer on the surface of complex-shaped substrates. However, the sprayed ceramic coating has a layered stacked structure, and the coating inevitably contains lamellar interfaces, pores, and longitudinal microcracks.

[0003] In related technologies, Al2O3 and MgAl2O4 are commonly used in thermal spraying insulating coatings. However, Al2O3 coatings are prone to forming interconnected cracks after high-temperature thermal cycling due to phase transformation, thermal expansion mismatch, and interlayer defects. MgAl2O4 coatings have good electrical insulation stability, but when used alone as a high-temperature coating for the hot end of aero-engines, its resistance to sintering and long-term phase stability still need improvement. Lanthanum magnesium hexaaluminate has good high-temperature phase stability and resistance to sintering, but when used as a single-phase powder in plasma spraying, it is difficult to form a continuous insulating closed phase at the lamellar interface after the molten droplets spread. During high-temperature insulation, insulation attenuation may still occur due to local sintering shrinkage and interconnected interlayer pores.

[0004] Furthermore, rare earth composite oxides such as rare earth tantalates and rare earth zircons exhibit good high-temperature phase stability, anti-sintering ability, or thermal expansion regulation in the field of thermal barrier coatings. Titanium-aluminum layered carbides and modified silicon nitride micropowders can fill microcracks through oxidation products or glassy phases in the field of high-temperature self-healing materials. Related technologies typically use these materials as the main material for thermal barrier coatings or the main reinforcing phase for self-healing structural ceramics, without localizing them into the same sprayed powder particles according to the functional differences between core and wall materials, nor utilizing their correspondence with the interfaces of insulating coating layers to address the problem of insulation degradation after high temperatures.

[0005] Regarding the aforementioned technologies, the inventors believe that currently available ceramic powder coatings are prone to problems such as increased interconnected porosity, increased thickness shrinkage, and disruption of the insulating pathways at the lamellar interfaces after exposure to high temperatures. Therefore, a ceramic powder capable of simultaneously depositing a high-temperature stable framework phase and a lamellar interface sealing phase within the same molten droplet is needed. Summary of the Invention

[0006] In order to improve the problem that non-core or mechanically mixed ceramic spray powders, after being kept at 1100℃ for high temperature, have increased interconnected porosity, increased coating thickness shrinkage, and decreased insulation retention due to asynchronous phase distribution, this application provides a ceramic powder and a dense insulating coating.

[0007] In a first aspect, this application provides a ceramic powder, which adopts the following technical solution: A ceramic powder comprising core-shell structured particles; the core-shell structured particles comprising a core material and a wall material covering the outer side of the core material; the core material comprising a lanthanum magnesium hexaaluminate phase and a framework regulating phase; the wall material comprising a magnesium aluminum spinel phase, a free α-Al2O3 phase, and a filler component, the filler component comprising MAX phase micro powder and / or modified silicon nitride micro powder; the ceramic powder comprising, by weight percentage, the following components: 55-68 wt% lanthanum magnesium hexaaluminate phase, 1-6 wt% framework regulating phase, 12-24 wt% magnesium aluminum spinel phase, 13-17 wt% free α-Al2O3 phase, with the balance made up to 100 wt% by the filler component; the D50 particle size of the ceramic powder is 25-55 μm.

[0008] By adopting the above technical solution, this application sets the lanthanum magnesium hexaaluminate phase and the framework regulating phase as the core material, and sets the magnesium aluminum spinel phase, free α-Al2O3 phase and filler components as the wall material. The D50 particle size of the ceramic powder is limited to 25-55μm, which can take into account both powder feeding stability and coating forming effect. It can avoid the failure of the core-shell structure due to excessive melting caused by the particle size being too small, and at the same time, it can avoid the insufficient melting of the particles caused by the particle size being too large.

[0009] In the core material, the lanthanum magnesium hexaaluminate phase exhibits good high-temperature phase stability and anti-sintering properties, providing a stable distribution carrier for the framework regulating phase during high-temperature insulation. The framework regulating phase, in turn, modulates the thermal expansion behavior of the core material, mitigating stress concentration caused by local thermal mismatch and non-uniform shrinkage of the lanthanum magnesium hexaaluminate phase. The spatial constraint effect of the lanthanum magnesium hexaaluminate phase on the framework regulating phase helps maintain its regulating effect, while the buffering effect of the framework regulating phase on local thermal stress helps maintain the integrity of the lanthanum magnesium hexaaluminate framework. The synergy of these two phases further inhibits high-temperature sintering shrinkage and high-temperature insulation damage in the core material, thereby improving its structural stability during high-temperature insulation. In the wall material, the magnesium aluminum spinel phase mitigates thermal stress at the lamellar interfaces and inhibits the generation and opening of microcracks, creating favorable conditions for the filler components to fill small pores and microcracks. The filler components' filling of pores and microcracks reduces stress concentration at crack tips, decreasing the likelihood of further crack propagation. The free α-Al₂O₃ phase and the magnesium aluminum spinel phase synergistically maintain the insulating properties of the ceramic powder. Through the above interactions, the wall material can simultaneously buffer interfacial stress, seal defects, and maintain insulation performance, thereby improving the bonding state between layers and reducing the interconnection of pores and microcracks.

[0010] Because the core and wall materials are bonded within the same particle in a core-shell structure, both components can enter the plasma flow simultaneously and deposit uniformly. The core material's suppression of lamellar sintering shrinkage reduces the tensile and shear deformation experienced by the wall material, allowing the interface buffering and defect sealing functions of the wall material to be effectively utilized. Conversely, the wall material's suppression of interfacial stress and interconnected defects reduces the concentration of local stress in the core material, thus preventing damage to the core skeleton due to the expansion of interface defects. Under the combined effects of the aforementioned component content, core-shell structure, and particle size range, a synergistic relationship is formed within the core material, the wall material, and between the core and wall materials. The resulting coating is less prone to significant sintering shrinkage and defect growth during high-temperature insulation at 1100℃, which helps reduce the coating's interconnected porosity and thickness variation rate, and improves the coating's insulation retention capacity after high-temperature insulation at 1100℃.

[0011] Preferably, the framework conditioning phase is a rare earth tantalate phase and / or a rare earth zirconate phase.

[0012] By adopting the above technical solution, this application sets rare earth tantalate phase and / or rare earth zirconate phase as one of the core material components of core-shell structured particles. This allows them to not directly undertake the function of sealing the lamellar interfaces, but rather to act as a framework regulating phase, forming a high-temperature stable framework together with the lanthanum magnesium hexaaluminate phase. The rare earth tantalate phase can regulate the thermal expansion behavior of the core material and reduce high-temperature sintering shrinkage; the rare earth zirconate phase has good anti-sintering and grain boundary pinning effects, which can inhibit grain growth and framework shrinkage of the core material after high-temperature heat treatment. When the two are combined, the rare earth zirconate phase provides an anti-sintering boundary for the rare earth tantalate phase, and the rare earth tantalate phase provides a high-temperature stress buffer for the rare earth zirconate phase. This makes the core material framework shrinkage and lamellar interface sealing more synchronized, improving the core material's ability to suppress high-temperature non-uniform shrinkage, which is beneficial for reducing the coating thickness change rate and improving the insulation retention rate after high temperature.

[0013] Preferably, the rare earth tantalate phase is at least one of YbTaO4, LuTaO4, and YTaO4; and the rare earth zirconate phase is at least one of Y2Zr2O7, Gd2Zr2O7, and La2Zr2O7.

[0014] By adopting the above technical solution, this application has selected specific types of rare earth tantalates and rare earth zircons, which is beneficial to reduce the coating thickness change rate and improve the insulation retention rate after high temperature.

[0015] Preferably, the filler component is MAX phase micro powder and / or modified silicon nitride micro powder; the general formula of the MAX phase is M n+1 AX n Where M is a transition metal element, A is a group IIIA or IVA main group element, X is carbon and / or nitrogen, and n=1, 2, 3.

[0016] By adopting the above technical solution, this application localizes MAX phase micropowder and / or modified silicon nitride micropowder in the wall material. MAX phase micropowder can form solid oxidation products such as Al2O3 and TiO2 under high-temperature oxidation conditions; the modified silicon nitride micropowder maintains a discrete distribution after spraying. When microcracks generated during high-temperature insulation and cooling expose silicon nitride to an oxygen-containing environment, the silicon nitride gradually oxidizes near the cracks to form a SiO2-based closed phase. The oxidation products of titanium-aluminum layered carbides can provide an attachment framework for the SiO2-based closed phase, which can then penetrate fine cracks and interconnected channels that are difficult for solid oxides to fill, improving the sealing effect of microcracks and interconnected channels at the layer interface. This is beneficial for reducing the porosity of the coating and improving the insulation retention rate after high temperatures.

[0017] Preferably, the MAX phase micro powder is at least one of Ti2AlC and Ti3AlC2; the modified silicon nitride micro powder includes silicon nitride micro powder and an alumina modification layer coated on the surface of the silicon nitride micro powder.

[0018] By adopting the above technical solution, this application has optimized the specific type of MAX phase micropowder and the compositional details of the modified silicon nitride micropowder. The alumina-modified layer, as a surface modification structure of the modified silicon nitride micropowder, should be included in the weight of the modified silicon nitride micropowder and calculated as part of the filler component, but not in the weight of the free α-Al₂O₃ phase. This avoids redundant measurement of the alumina-modified layer and the free α-Al₂O₃ phase, ensuring a clear boundary between the filler component and the free α-Al₂O₃ phase content; these filler components are beneficial for reducing the interconnected porosity of the coating and improving the insulation retention rate after high temperature.

[0019] Preferably, the thickness of the wall material is 1.5-3 μm.

[0020] By adopting the above technical solution, the wall material thickness determines whether the insulating sealing phase and filler components can preferentially cover the lamellar interface in the initial stage of spraying. When the wall material is too thin, the insulating sealing phase is insufficient to continuously cover the lamellar boundary on a large scale; when the wall material is too thick, the overall melting behavior of the particles tends to be similar to that of spinel-alumina powder, and the high-temperature stabilizing skeleton effect of the core material is weakened. The preferred wall material thickness of this application improves the continuous coverage and defect sealing effect of the wall material at the lamellar interface, enabling the coating to achieve a combination of low interconnected porosity, low coating thickness variation rate, and high insulation retention rate after high temperature.

[0021] Preferably, the ceramic powder is prepared according to the following steps: preparing a lanthanum magnesium hexaaluminate pre-synthesized powder containing a framework regulating phase; adding the pre-synthesized powder to a coating slurry containing pre-formed magnesium aluminum spinel micro powder, free α-Al2O3 powder and filler components and grinding and dispersing it to obtain a composite slurry; centrifugally spray-granulating the composite slurry; and sieving the granulated powder after heat treatment in an inert atmosphere to obtain ceramic powder.

[0022] By adopting the above technical solution, this application optimizes the preparation steps of ceramic powder. In the above steps, lanthanum magnesium hexaaluminate pre-synthesized powder constitutes the core particles in the composite slurry, and the D50 of the pre-synthesized powder is smaller than the D50 of the finished ceramic powder. After the composite slurry is centrifuged and atomized to form droplets, water evaporates from the surface of the droplets to the outside, the pre-synthesized powder is retained inside the droplets, and the wall material components are enriched and solidified on the outside; subsequently, after argon heat treatment, core-shell structured particles can be obtained.

[0023] Preferably, the magnesium-aluminum spinel phase is provided by pre-formed magnesium-aluminum spinel micro powder added to the coating slurry.

[0024] By adopting the above technical solution, the preformed magnesium aluminum spinel phase is used in the wall material to buffer the thermal stress at the interface of the sheet layer, which improves the effect of the wall material in buffering the thermal stress at the interface of the sheet layer and maintaining the insulation path. This is beneficial to reducing the coating connectivity porosity and coating thickness variation rate, and improving the insulation retention rate after high temperature.

[0025] Secondly, this application provides a dense insulating coating, employing the following technical solution: A dense insulating coating includes an adhesive layer and a ceramic surface layer sequentially disposed on the surface of a substrate; the adhesive layer is a NiCrAlY layer; the ceramic surface layer is formed by plasma spraying of the aforementioned ceramic powder, and the thickness of the ceramic surface layer is 300-500 μm.

[0026] By adopting the above technical solution, the NiCrAlY bonding layer serves to buffer the difference in thermal expansion between the substrate and the ceramic surface layer, while the ceramic surface layer undertakes the function of high-temperature insulation. This thickness range is not simply an optimization of coating thickness, but rather ensures that the lamellar structure formed by the deposition of core-shell structured particles is continuously connected in the thickness direction: when the ceramic surface layer is too thin, the number of repeated stackings of the central skeleton and the wall material closed structure is insufficient; when the ceramic surface layer is too thick, the accumulation of interlayer stress generated during high-temperature insulation and cooling processes increases.

[0027] Preferably, the ceramic surface layer includes a dense transition zone near the adhesive layer and a main body zone located outside the dense transition zone. The spraying current of the dense transition zone is greater than that of the main body zone, and the difference between the two is 100-150A.

[0028] By adopting the above technical solution, this application uses a higher spraying current in the dense transition zone to fully melt the core-shell structure particle wall material and preferentially seal the pores near the adhesive layer; a lower spraying current is used in the main body zone to prevent excessive melting of the lanthanum magnesium hexaaluminate phase and the framework conditioning phase in the core material. These two spraying parameters enable a smooth transition between the adhesive layer and the ceramic surface layer, helping to extend the service life of the ceramic surface layer.

[0029] In summary, this application has the following beneficial effects: 1. In this application, the lanthanum magnesium hexaaluminate phase and the framework regulating phase are used as the core material, and the magnesium aluminum spinel phase, the free α-Al2O3 phase and the filler components are used as the wall material. The core material and the wall material are combined into core-shell structured particles, so that the core material, the wall material and the core material and the wall material form a mutually synergistic relationship. This is beneficial to reduce the interconnected porosity and coating thickness change rate after high temperature insulation at 1100℃, and improve the insulation retention rate after high temperature.

[0030] 2. This application further optimizes the specific types of the skeleton conditioning phase and filler components, so that the effect of the core material in suppressing non-uniform shrinkage at high temperature and the effect of the wall material in sealing microcracks and connecting channels at the layer interface can be further enhanced. This is beneficial to reducing the coating thickness change rate and the interconnected porosity, and improving the insulation retention rate after high temperature.

[0031] 3. By setting the wall material thickness, ceramic surface layer thickness, and spraying current difference between the dense transition zone and the main body zone, this application enables the core material-wall material structure to continuously function at the layer interface and in the coating thickness direction, which is beneficial to make the coating have a low interconnected porosity, a low coating thickness change rate, and a high insulation retention rate after high temperature. Detailed Implementation

[0032] The present application will be further described in detail below with reference to examples, preparation examples and comparative examples. For ease of traceability, the raw materials used in the examples of this application are listed as follows: La2O3 is a powder with a purity of not less than 99.9% and a D50 of 1.5-3.0 μm. Before use, it is kept at 900℃ for 2 hours and cooled in a desiccator for later use; MgO is a powder with a purity of not less than 99.9% and a D50 of 1.0-2.5 μm. Before use, it is kept at 900℃ for 2 hours and cooled in a desiccator for later use; α-Al2O3 is a powder with a purity of not less than 99.9% and a D50 of 0.20-0.40 μm, with a D90 not greater than 0.50 μm; YbTaO4, L Both uTaO4 and YTaO4 are powders with a purity of not less than 99.0% and a D50 of 0.7-1.1 μm. The aforementioned rare earth tantalate powders can be commercially available powders meeting these purity and particle size requirements, or they can be obtained by ball milling, drying, pre-calcining at 1400-1550℃ for 3-6 hours, crushing, and grading of the corresponding rare earth oxides and Ta2O5 in a stoichiometric ratio. Y2Zr2O7, Gd2Zr2O7, and La2Zr2O7 are all powders with a purity of not less than 99.0% and a D50 of 0.7-1.1 μm. The aforementioned rare earth zirconate powders can be obtained by ball milling, drying, pre-calcining at 1400-1550℃ for 3-6 hours, crushing, and grading of the corresponding rare earth oxides and Ta2O5 in a stoichiometric ratio. Commercially available powders with the required particle size can also be obtained by ball milling, drying, pre-calcining at 1400-1550℃ for 3-6 hours, crushing, and classifying the corresponding rare earth oxides and ZrO2 in stoichiometric ratio; Ti2AlC is a powder with a purity of not less than 98%, and after air jet milling and classification, the D50 is 0.25-0.35μm, and the D90 is not greater than 0.45μm; Ti3AlC2 is a powder with a purity of not less than 98%, and after air jet milling and classification, the D50 is 0.30-0.40μm, and the D90 is not greater than 0.50μm; ammonium polyacrylate is a dispersant for ceramic slurries, and ammonium citrate is an analytical... The materials used are: pure NiCrAlY powder (for thermal spraying, with a particle size range of 45-106 μm); MgAl2O4 preform powder (with a purity of not less than 99.0%, and a D50 of 0.25-0.35 μm and a D90 of not more than 0.45 μm after air jet milling and classification); spherical silicon nitride micropowder (with a purity of not less than 99.0%, and a D50 of 0.25-0.35 μm and a D90 of not more than 0.45 μm after air jet milling and classification); aluminum isopropoxide (with a purity of not less than 98%); and isopropanol, polyvinylpyrrolidone, and nitric acid (all conventional chemical reagents). The sources of the above raw materials are not limited; equivalent raw materials meeting the same phase composition, purity, and particle size requirements can be used.

[0033] Preparation example of modified silicon nitride micro powder The following explanation uses Preparation Example 1 as an example.

[0034] Preparation Example 1 In this preparation example, the modified silicon nitride micro powder was prepared according to the following method: (1) Weigh 20.00g of spherical silicon nitride micro powder, 120g of isopropanol and 0.20g of polyvinylpyrrolidone, place them in a polytetrafluoroethylene container, and stir at 400rpm for 30min to obtain a silicon nitride dispersion; the spherical silicon nitride micro powder has a D50 of 0.30μm, a D90 of no more than 0.45μm, and a purity of no less than 99%.

[0035] (2) Add 40g of isopropanol solution containing 7.00g of aluminum isopropoxide to the silicon nitride dispersion obtained in step (1), then add 1.50g of deionized water, adjust the pH to 4.0 with nitric acid, stir at 60℃ for 4h, so that aluminum isopropoxide hydrolyzes on the surface of silicon nitride micro powder and forms a continuous aluminum hydrated oxide coating layer, and obtain the coated precursor slurry.

[0036] (3) The coated precursor slurry obtained in step (2) is centrifuged and washed three times with isopropanol. After drying at 80°C for 8 hours, it is heated to 750°C at 3°C / min and kept at that temperature for 2 hours in an argon atmosphere to dehydrate the aluminum-containing hydrated oxide and form an alumina modified layer, thereby obtaining modified silicon nitride micro powder. The D50 of the obtained modified silicon nitride micro powder is 0.35μm, and the D90 is not greater than 0.50μm. The alumina modified layer accounts for 8wt% of the weight of the modified silicon nitride micro powder and is included as part of the filler component, excluding the free α-Al2O3 phase.

[0037] Example Examples 1-5

[0038] The following description uses Example 1 as an example.

[0039] In this embodiment, the ceramic powder comprises core-shell structured particles; the core-shell structured particles comprise a core material and a wall material covering the outside of the core material. The core material comprises a lanthanum magnesium hexaaluminate phase and a framework regulating phase, and the wall material comprises a magnesium aluminum spinel phase, a free α-Al₂O₃ phase, and a filler component. Based on the total weight of the ceramic powder, the ceramic powder comprises 55 wt% lanthanum magnesium hexaaluminate phase, 1 wt% framework regulating phase, 24 wt% magnesium aluminum spinel phase, 17 wt% free α-Al₂O₃ phase, and 3 wt% filler component; the D50 particle size of the ceramic powder is 25 μm.

[0040] In this embodiment, the framework conditioning phase is YbTaO4, the filler component is Ti2AlC, the wall material thickness is 1.2μm, and the magnesium aluminum spinel phase is provided by pre-formed MgAl2O4 micro powder added to the coating slurry.

[0041] In this embodiment, the contents of each component listed in Table 1 are all weight percentages of each component in the finished ceramic powder. During preparation, the pre-synthesized lanthanum magnesium hexaaluminate powder containing the framework regulating phase, the pre-formed MgAl2O4 micro powder, the free α-Al2O3 powder, and the filler components were weighed according to the composition of the finished product listed in Table 1. In Examples 1 to 5, YbTaO4 was used as the framework regulating phase and Ti2AlC was used as the filler component, with the filler component being added to a balance of 100 wt%.

[0042] This embodiment provides a method for preparing ceramic powder, including preparing a core material pre-synthesized powder, formulating a coating slurry containing pre-formed MgAl2O4 micro powder and filler components, centrifugal spray granulation, and argon heat treatment, specifically including the following steps: (1) Weigh 11.75g ​​of La2O3, 2.91g of MgO, 40.45g of α-Al2O3 and 1.00g of YbTaO4, place them in a nylon ball mill jar, add 300g of zirconia balls and 120g of deionized water, and ball mill at 300rpm for 12h. Dry the ball-milled slurry at 110℃ for 10h, pass it through an 80-mesh sieve and place it in an alumina crucible, pre-synthesize it at 1600℃ for 4h with a heating rate of 5℃ / min, cool it in the furnace, crush it, wet ball mill it for 8h and classify it to obtain a lanthanum magnesium hexaaluminate pre-synthesized powder containing rare earth tantalate phase with a D50 of 10μm and a D90 of no more than 20μm. The pre-synthesized powder includes 55.00g of lanthanum magnesium hexaaluminate phase and 1.00g of YbTaO4 according to the target phase.

[0043] (2) Weigh 56.00g of the pre-synthesized powder, 24.00g of the pre-formed MgAl2O4 micro powder, 17.00g of the free α-Al2O3 powder, 3.00g of Ti2AlC, 0.30g of ammonium polyacrylate and 0.15g of ammonium citrate obtained in step (1), add them to 150g of deionized water, adjust the pH to 9.2 with ammonia water, stir at 500rpm for 40min and then sand mill for 2h to obtain a composite slurry; wherein the D90 of the pre-formed MgAl2O4 micro powder, the free α-Al2O3 powder and Ti2AlC are all not greater than 0.50μm, and can be used as wall material components to adhere to the surface of the pre-synthesized powder; the pre-formed MgAl2O4 micro powder directly provides the magnesium aluminum spinel phase in the finished ceramic powder, and the free α-Al2O3 powder directly provides the free α-Al2O3 phase in the finished ceramic powder.

[0044] (3) The composite slurry obtained in step (2) is centrifuged and sprayed to granulate. The inlet air temperature is 220℃, the outlet air temperature is 95℃, and the atomizing disc speed is 12000rpm. The granulated powder is heat-treated at 950℃ for 2h in argon atmosphere. The 15-45μm fraction is sieved to obtain ceramic powder with D50 of 25μm.

[0045] This embodiment also provides a method for preparing a dense insulating coating, comprising the following steps: sandblasting a nickel-based alloy substrate to a surface roughness Ra of 4.5 μm, and then spraying a 60 μm thick NiCrAlY bonding layer using Metco 461NS powder; subsequently, spraying a 260 μm thick ceramic surface layer using the ceramic powder obtained in this embodiment. The ceramic surface layer includes a dense transition region near the bonding layer and a main body region located outside the dense transition region, wherein the spraying current for the dense transition region is 600 A, the spraying current for the main body region is 510 A, the spraying voltage is 65 V, the spraying distance is 105 mm, and the powder feed rate is 28 g / min.

[0046] The difference between Examples 2-5 and Example 1 lies in the content of the framework regulating phase, the amount of filler component, the corresponding total formulation, and the D50 particle size in the core-shell structured particles. The framework regulating phase in Examples 2-5 is also YbTaO4, and the filler component is also Ti2AlC. Except for the formulation, D50 particle size, and corresponding raw material input adjustments listed in Table 1, the raw material types, basic preparation processes, and spraying process parameters used in Examples 2-5 are the same as in Example 1.

[0047] In Table 1, the D50 particle size was obtained by adjusting the spray granulation conditions and sieving grades. The core material pre-synthesized powder used in Examples 1-5 was controlled to have a D50 of 8-15 μm and a D90 of no more than 22 μm. The D90 of the wall material powder used was no more than 0.50 μm.

[0048] Table 1. Specific formulations and D50 particle sizes for Examples 1 to 5.

[0049] Examples 6-12

[0050] As shown in Table 2, the difference between Examples 6-12 and Example 3 is that the specific types of the skeleton adjustment phase in the core material are different; except for the specific types of skeleton adjustment phases and the compounding mass ratio listed in Table 2, the other conditions are the same as those in Example 3.

[0051] Table 2. Types of skeleton-modifying phases and their mass ratios

[0052] Examples 13-17

[0053] As shown in Table 3, the difference between Examples 13-17 and Example 9 is that the specific types of filler components are different; the modified silicon nitride micro powder used in Examples 14, 16 and 17 are all modified silicon nitride micro powders obtained in Preparation Example 1. The modified silicon nitride micro powder is included in the filler components according to the total weight, and the free α-Al2O3 phase is not included in the alumina modification layer; except for the specific types of filler components and the compounding mass ratio listed in Table 3, the other conditions are the same as in Example 9.

[0054] Table 3. Types of filler components and their mass ratios

[0055] Examples 18-22

[0056] As shown in Table 4, the difference between Examples 18-22 and Example 16 is that the wall material thickness is different; except for the wall material thickness listed in Table 4, the other conditions are the same as those in Example 16.

[0057] Table 4 Wall Material Thickness

[0058] Examples 23-27

[0059] As shown in Table 5, the difference between Examples 23-27 and Example 20 is that the thickness of the ceramic surface layer is different; except for the ceramic surface layer thickness listed in Table 5, the other conditions are the same as those in Example 20.

[0060] Table 5 Ceramic Surface Layer Thickness

[0061] Examples 28-32

[0062] As shown in Table 6, the difference between Examples 28-32 and Example 25 is that the spraying current difference between the dense transition zone and the main body zone is different. Specifically, the spraying current of the main body zone remains unchanged, while only the spraying current of the dense zone is changed. Except for the spraying current difference listed in Table 6, the other conditions are the same as in Example 25.

[0063] Table 6. Spraying current difference between the dense transition zone and the main body zone

[0064] Comparative Example Comparative Example 1 The difference between Comparative Example 1 and Example 1 is that core-shell structured particles are not prepared. Instead, lanthanum magnesium hexaaluminate powder, YbTaO4 powder, magnesium aluminum spinel powder, free α-Al2O3 powder and Ti2AlC powder are mechanically mixed according to the five-phase conversion formula of Example 1 and used as a spraying powder.

[0065] Comparative Example 2 The difference between Comparative Example 2 and Example 1 is that the core material does not contain the lanthanum magnesium hexaaluminate phase, but is replaced by magnesium aluminum spinel phase and free α-Al2O3 phase.

[0066] Comparative Example 3 The difference between Comparative Example 3 and Example 1 is that the core material does not contain the framework adjustment phase, and the corresponding weight of the framework adjustment phase is made up by the lanthanum magnesium hexaaluminate phase.

[0067] Comparative Example 4 The difference between Comparative Example 4 and Example 1 is that the wall material does not contain the magnesium aluminum spinel phase, and the corresponding weight of the magnesium aluminum spinel phase is made up by the free α-Al2O3 phase.

[0068] Comparative Example 5 The difference between Comparative Example 5 and Example 1 is that the wall material does not contain the free α-Al2O3 phase, and the corresponding weight of the free α-Al2O3 is made up by the magnesium aluminum spinel phase.

[0069] Comparative Example 6 The difference between Comparative Example 6 and Example 1 is that the wall material does not contain any filler components, and the corresponding weight of the filler components is made up by the magnesium aluminum spinel phase.

[0070] Comparative Example 7 The difference between Comparative Example 7 and Example 1 is that the D50 particle size of the ceramic powder is 12 μm.

[0071] Comparative Example 8 The difference between Comparative Example 8 and Example 1 is that the D50 particle size of the ceramic powder is 72 μm.

[0072] Performance testing methods Performance testing: At least 3 parallel test specimens shall be prepared for each group of samples; unless otherwise specified, the test results shall be the average value of the parallel test specimens.

[0073] I. Porosity Test: Sample sections were cut perpendicular to the coating surface. After cold mounting with epoxy resin, the sections were sequentially smoothed with 400#, 800#, 1200#, and 2000# SiC sandpaper, and then polished with 0.05μm alumina polishing liquid until no obvious scratches remained. The ceramic surface section was observed using a scanning electron microscope in backscatter mode or a metallographic microscope. At least 10 fields of view were randomly selected along the coating length on each sample, with a magnification of 500x. ImageJ or equivalent image analysis software was used to threshold-segment the pore grayscale region and the ceramic matrix region. Porosity was calculated as the percentage of pore area to the total cross-sectional area of ​​the ceramic surface.

[0074] II. Connected Porosity Test: The sample was placed in low-viscosity epoxy resin containing fluorescent dye, and a vacuum was applied below -0.09 MPa for 30 minutes. Then, the pressure was restored to normal and the sample was impregnated for another 2 hours to allow the resin to penetrate into pores and cracks connected to the outside environment. After the resin cured, a cross-section was prepared using the same method as the porosity test. The ceramic surface cross-section was observed using a fluorescence microscope. At least 10 fields of view were randomly selected on each sample, and the area of ​​the fluorescent resin-filled region was counted. The connected porosity was calculated as the percentage of the fluorescent resin-filled region to the total area of ​​the ceramic surface cross-section. Closed pores that were not connected to the outside environment and were not filled with fluorescent resin were not included in the connected porosity.

[0075] III. Coating Thickness Change Rate Test: The thickness of the ceramic surface layer was measured on the cross-section of the sample before and after high-temperature insulation. For each sample, at least 10 measuring points were selected at equal intervals along its length. The distance from the NiCrAlY adhesive layer / ceramic surface layer interface to the outer surface of the ceramic surface layer was measured, and the average value was taken as the ceramic surface layer thickness of the sample. The coating thickness change rate was calculated using the formula: (absolute value of (average thickness before insulation - average thickness after insulation) / average thickness before insulation) × 100%.

[0076] IV. High-Temperature Service Test: The sample was placed on an alumina firing plate and placed in an air atmosphere muffle furnace. The temperature was increased to 1100℃ at a rate of 5℃ / min, and held at 1100±5℃ for 100 hours. After cooling in the furnace, the sample was removed below 100℃. After cooling, the coating surface was first observed using a stereomicroscope, and then a cross-section was prepared and observed using a metallographic microscope or scanning electron microscope. If a crack penetrated the thickness of the ceramic surface layer or if flaking was visible to the naked eye, it was recorded as a through crack or flaking.

[0077] V. Insulation Retention Rate Test After High Temperature: The volume resistivity of the samples was tested before and after the high-temperature service test. Before the test, a 10mm diameter silver or platinum electrode was prepared on the ceramic surface, with the metal substrate used as the back electrode. The sample edges were sealed with high-temperature resistant insulating edge sealant and then dried to reduce edge leakage. During the test, a 500V DC voltage was applied at room temperature and relative humidity not exceeding 60%. After stabilizing for 60 seconds, the resistance R was read. The volume resistivity was calculated using ρv = R × A / d, where A is the surface electrode area and d is the ceramic surface thickness. The insulation retention rate after high temperature was calculated using ρv. (保温后) / ρv (保温前) Calculated by multiplying by 100%.

[0078] Table 7

[0079]

[0080] Based on Examples 1 to 5 and Comparative Examples 1 to 8, and in conjunction with Table 7, it can be seen that in Examples 1 to 5, with a full coverage of 1-6 wt% of the skeleton adjustment phase, no through-cracks appeared after heat preservation at 1100℃, the interconnected porosity was 0.78-1.28%, the coating thickness change rate was 1.46-2.05%, and the insulation retention rate after high temperature was 82-92%. In contrast, Comparative Examples 1 to 8, due to the lack of core-shell structure particles, core material lanthanum magnesium hexaaluminate phase, skeleton adjustment phase, wall material magnesium aluminum spinel phase, free α-Al2O3 phase, filler components, or D50 particle size conditions, had interconnected porosity increased to 2.45-4.20%, coating thickness change rate increased to 3.38-5.35%, and insulation retention rate after high temperature decreased to 31-57%. The above data shows that the component content, core-shell structure, and D50 particle size range can create a synergistic effect within the core material, the wall material, and between the core material and the wall material. This reduces sintering shrinkage and defect growth after high-temperature insulation at 1100℃, which helps to reduce the interconnected porosity and coating thickness variation rate of the coating, and improves the insulation retention rate after high temperature.

[0081] Further analysis of Comparative Example 1 reveals that even when the calculated formulations of each component are the same as in Example 1, as long as mechanical mixing is used and core-shell structured particles are not formed, the interconnected porosity and coating thickness change rate still increase significantly after high temperature. This is because different components in the mechanically mixed powder cannot enter the plasma flow simultaneously with the same particles and deposit uniformly. The core material and wall material are difficult to interact with, resulting in a weakened sealing effect of the wall material on interfacial defects and a weakened inhibitory effect of the core material on sintering shrinkage. Therefore, the insulation retention rate measured in Comparative Example 1 after high temperature is significantly lower than that in Example 1.

[0082] Further analysis of Comparative Examples 2 to 6 reveals that when the core material lacks the lanthanum-magnesium hexaaluminate phase, the framework regulating phase lacks a stable distribution carrier; when the core material lacks the framework regulating phase, the local thermal mismatch and non-uniform shrinkage of the lanthanum-magnesium hexaaluminate phase are difficult to mitigate; when the wall material lacks the magnesium-aluminum spinel phase, the thermal stress at the lamellar interface is difficult to buffer; when the wall material lacks the free α-Al₂O₃ phase, the interfacial insulation continuity is insufficient; and when the wall material lacks filler components, there is a lack of local filler sources near the microcracks. Therefore, Comparative Examples 2 to 6 all exhibit problems such as increased interconnected porosity, increased coating thickness variation rate, and decreased insulation retention rate after high temperature, indicating that the components in each region need to be used in combination, rather than simply in parallel.

[0083] Further analysis of Comparative Examples 7 and 8 reveals that when the D50 particle size is too small, the powder is prone to over-melting or deviating from the deposition area, making it difficult to stably retain the core-shell structure. When the D50 particle size is too large, the particles do not melt sufficiently, resulting in an increase in unmelted particles and unspread pores in the coating. Therefore, the test data of Comparative Examples 7 and 8 are worse than those of Example 1, indicating that the D50 particle size range needs to be used in conjunction with the core-shell spatial structure to balance powder delivery stability and coating formation effect.

[0084] As can be seen from Examples 6 to 12 and Table 7, based on the basic structures of Examples 1 to 5, when the skeleton adjustment phase is replaced with LuTaO4, YTaO4, Y2Zr2O7, YTaO4 / Y2Zr2O7=1:1, YbTaO4 / Y2Zr2O7=1:1, Gd2Zr2O7 or La2Zr2O7, the coating porosity is 1.45-2.10%, the interconnected porosity is 0.44-0.75%, the coating thickness variation rate is 0.98-1.40%, and the insulation retention rate after high temperature is 89.5-95.0%, all of which remain at a good level. In Example 9, YTaO4 / Y2Zr2O7 = 1:1 was used, resulting in lower interconnected porosity and coating thickness variation rate, and higher insulation retention rate after high temperature. When Gd2Zr2O7 and La2Zr2O7 were used alone as rare earth zirconate phases to regulate the skeleton, they could also form a high-temperature stable skeleton together with the lanthanum magnesium hexaaluminate phase, suppressing the non-uniform shrinkage of the core material after high-temperature insulation. When rare earth tantalate phase and rare earth zirconate phase were compounded, the rare earth zirconate phase provided an anti-sintering boundary for the rare earth tantalate phase, and the rare earth tantalate phase provided a high-temperature stress buffer for the rare earth zirconate phase. The two together improved the effect of the core material in suppressing non-uniform shrinkage at high temperature, which is beneficial to reducing the coating thickness variation rate and improving the insulation retention rate after high temperature.

[0085] Referring to Examples 13 to 17 and Table 7, it can be seen that, based on the framework adjustment phase setting of Example 9, when the filler component is changed to Ti3AlC2, the modified silicon nitride micropowder obtained in Preparation Example 1, Ti2AlC / Ti3AlC2 = 1:1, Ti2AlC / modified silicon nitride micropowder obtained in Preparation Example 1 = 1:1, or Ti3AlC2 / modified silicon nitride micropowder obtained in Preparation Example 1 = 1:1, the coating interconnected porosity is 0.40-0.55%, the coating thickness change rate is 0.84-1.06%, and the insulation retention rate after high temperature is 93.6-95.8%, all remaining at a good level. Among them, Example 16, using Ti2AlC / modified silicon nitride micropowder obtained in Preparation Example 1 = 1:1, has a lower interconnected porosity and coating thickness change rate, and a higher insulation retention rate after high temperature. The above data shows that both MAX phase micro powder and modified silicon nitride micro powder can be combined with wall material oxide sealing phase. When MAX phase micro powder is combined with modified silicon nitride micro powder obtained in Preparation Example 1, the solid oxide formed by the oxidation of MAX phase provides an adhesion interface for SiO2-based glass phase. The SiO2-based glass phase also makes up for the problem of insufficient fluidity of oxidation products of titanium-aluminum layered carbides, which further improves the sealing effect of microcracks and interconnected channels at the layer interface, which is beneficial to reduce the interconnected porosity of the coating and improve the insulation retention rate after high temperature.

[0086] As can be seen from Examples 18 to 22, a good coating structure and high-temperature insulation retention effect can be obtained when the wall material thickness is between 1.5-3 μm. As the wall material thickness increases from 1.5 μm to 2.0 μm, the porosity decreases from 1.24% to 1.02%, the interconnected porosity decreases from 0.38% to 0.27%, the coating thickness change rate decreases from 0.79% to 0.60%, and the high-temperature insulation retention rate increases from 96.2% to 97.5%. When the wall material thickness continues to increase to 3.0 μm, the above performance declines. Therefore, the overall performance is better when the wall material thickness is 2.0 μm. This result indicates that within the above range, appropriately increasing the wall material thickness is beneficial to improving the continuous coverage of the lamellar interface and the defect sealing effect of the wall material; however, when the wall material is too thick, the balance between the overall melting behavior of the particles and the retention of the core skeleton is affected, and the improvement in overall performance is reduced.

[0087] As can be seen from Examples 23 to 27, a testable continuous insulating layer can be formed when the ceramic surface layer thickness is between 300 and 500 μm. As the ceramic surface layer thickness increases from 300 μm to 400 μm, the porosity decreases from 0.96% to 0.80%, the interconnected porosity decreases from 0.25% to 0.19%, the coating thickness change rate decreases from 0.56% to 0.44%, and the insulation retention rate after high temperature increases from 97.7% to 98.8%. When the ceramic surface layer thickness continues to increase to 500 μm, the above performance declines. Therefore, the overall performance is better when the ceramic surface layer thickness is 400 μm. This result indicates that within the above range, increasing the ceramic surface layer thickness is beneficial for the repeated stacking of core-shell structure powder deposition sheets in the thickness direction to form a continuous insulating barrier; however, when the thickness is too large, the interlayer stress accumulation during high-temperature insulation and cooling processes increases, and the improvement in overall performance decreases.

[0088] As can be seen from Examples 28 to 32, when the spraying current difference between the dense transition zone and the main body zone is between 100-150A, the ceramic surface layer can achieve both densification near the adhesive layer and skeletal preservation in the main body zone. As the spraying current difference increases from 100A to 130A, the porosity decreases from 0.77% to 0.67%, the interconnected porosity decreases from 0.18% to 0.14%, the coating thickness change rate decreases from 0.42% to 0.34%, and the insulation retention rate after high temperature increases from 99.0% to 99.5%. When the spraying current difference continues to increase to 150A, the above performance declines. Therefore, the overall performance is better when the current difference is 130A. The above settings are beneficial for reducing the interconnected porosity and coating thickness change rate after high temperature, and improving the insulation retention rate after high temperature; however, when the current difference is too large, the microstructural difference between the dense transition zone and the main body zone increases, and the improvement in overall performance decreases.

[0089] The above embodiments are merely explanations of this application and are not intended to limit it. After reading this specification, those skilled in the art can make modifications to the embodiments of this application without contributing any inventive step, but such modifications are protected by patent law as long as they fall within the scope of the claims of this application.

Claims

1. A ceramic powder, characterized in that, The ceramic powder is a core-shell structure particle; the core-shell structure particle includes a core material and a wall material covering the outside of the core material; the core material contains a lanthanum magnesium hexaaluminate phase and a framework regulating phase; the wall material contains a magnesium aluminum spinel phase, a free α-Al2O3 phase, and a filler component; the ceramic powder comprises the following components by weight percentage: 55-68 wt% lanthanum magnesium hexaaluminate phase, 1-6 wt% framework regulating phase, 12-24 wt% magnesium aluminum spinel phase, 13-17 wt% free α-Al2O3 phase, with the balance made up to 100 wt% by the filler component; the D50 particle size of the ceramic powder is 25-55 μm; the framework regulating phase is a rare earth tantalate phase and / or a rare earth zirconate phase; the filler component is MAX phase micro powder and / or modified silicon nitride micro powder; the general formula of the MAX phase is M n+1 AX n Where M is a transition metal element, A is a group IIIA or IVA main group element, X is carbon and / or nitrogen, and n=1, 2, 3.

2. The ceramic powder according to claim 1, characterized in that, The rare earth tantalate phase is at least one of YbTaO4, LuTaO4, and YTaO4; the rare earth zirconate phase is at least one of Y2Zr2O7, Gd2Zr2O7, and La2Zr2O7.

3. The ceramic powder according to claim 1, characterized in that, The MAX phase micro powder is at least one of Ti2AlC and Ti3AlC2; the modified silicon nitride micro powder includes silicon nitride micro powder and an alumina modification layer coated on the surface of the silicon nitride micro powder.

4. The ceramic powder according to claim 1, characterized in that, The wall material has a thickness of 1.5-3 μm.

5. The ceramic powder according to claim 1, characterized in that, The ceramic powder is prepared according to the following steps: a pre-synthesized powder of lanthanum magnesium hexaaluminate containing a framework regulating phase is prepared; the pre-synthesized powder is added to a coating slurry containing pre-formed magnesium aluminum spinel micro powder, free α-Al2O3 powder and filler components and ground and dispersed to obtain a composite slurry; the composite slurry is centrifugally spray-granulated; the granulated powder is heat-treated in an inert atmosphere and then sieved to obtain ceramic powder.

6. The ceramic powder according to claim 5, characterized in that, The magnesium-aluminum spinel phase is provided by pre-formed magnesium-aluminum spinel micro powder added to the coating slurry.

7. A dense insulating coating, characterized in that, It includes an adhesive layer and a ceramic surface layer sequentially disposed on the surface of a substrate; the adhesive layer is a NiCrAlY layer; the ceramic surface layer is formed by plasma spraying of the ceramic powder according to any one of claims 1-6, and the thickness of the ceramic surface layer is 300-500μm.

8. The dense insulating coating according to claim 7, characterized in that, The ceramic surface layer includes a dense transition zone near the adhesive layer and a main body zone located outside the dense transition zone. The spraying current of the dense transition zone is greater than that of the main body zone, and the difference between the two is 100-150A.

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